Densification, Tailored Microstructure, and Mechanical Properties of Selective Laser Melted Ti–6Al–4V Alloy via Annealing Heat Treatment

This work investigated the influence of process parameters on the densification, microstructure, and mechanical properties of a Ti–6Al–4V alloy printed by selective laser melting (SLM), followed by annealing heat treatment. In particular, the evolution mechanisms of the microstructure and mechanical properties of the printed alloy with respect to the annealing temperature near the β phase transition temperature were investigated. The process parameter optimization of SLM can lead to the densification of the printed Ti–6Al–4V alloy with a relative density of 99.51%, accompanied by an ultimate tensile strength of 1204 MPa and elongation of 7.8%. The results show that the microstructure can be tailored by altering the scanning speed and annealing temperature. The SLM-printed Ti–6Al–4V alloy contains epitaxial growth β columnar grains and internal acicular martensitic α′ grains, and the width of the β columnar grain decreases with an increase in the scanning speed. Comparatively, the printed alloy after annealing in the range of 750–1050 °C obtains the microstructure consisting of α + β dual phases. In particular, network and Widmanstätten structures are formed at the annealing temperatures of 850 °C and 1050 °C, respectively. The maximum elongation of 14% can be achieved at the annealing temperature of 950 °C, which was 79% higher than that of as-printed samples. Meanwhile, an ultimate tensile strength larger than 1000 MPa can be maintained, which still meets the application requirements of the forged Ti–6Al–4V alloy.


Introduction
Ti-6Al-4V alloy has been widely used in the aerospace, energy, biomedical, and automotive sectors [1,2] due to its high strength, low density, high fracture toughness, excellent corrosion resistance, and good biocompatibility [3]. Metal additive manufacturing (AM) has been advancing in the fabrication of geometrically complex metal products, typically including selective laser melting (SLM), directed energy deposition, metal binder jetting, and sheet lamination [4][5][6]. SLM has been widely applied to manufacture complex titanium parts with short lead time, great design freedom, and comparable product performance to forged counterparts [7], such as aircraft brackets [8], cervical fusion cages [9], bone implants [10], and partial denture clasps [11].
A considerable number of research works have been conducted on the fabrication of the Ti-6Al-4V alloy via SLM. Process parameters have great influence on the relative

Materials
An atomized Ti-6Al-4V alloy powder (AP&C company, Boisbriand, QC, Canada) with an average particle size of 33 µm was used. The material composition is shown in Table 1. The loose density of the powder was 2.45 g/cm 3 . Figure 1 shows the morphology of the powder and its particle size distribution. The powder particles were almost entirely spherical, and the particle size distribution was 15-45 µm.

SLM Process and Heat Treatment
A Dimetal-100 SLM equipment (Laseradd Technology Co., Ltd., Guangzhou, Guangdong, China) was utilized to print the Ti-6Al-4V alloy powder. The process optimization was conducted, and the parameter variables are shown in Table 2. The scanning strategy was bidirectional orthogonal scanning with a scanning starting angle of 145°. The cubic samples with dimensions of 8 mm × 8 mm × 8 mm and tensile samples were vertically printed, respectively, using various laser power, scanning speed, and hatch space values. All the samples were printed at the same position, i.e., central area of the substrate. Table 2. Process parameters of selective laser melting (SLM) for printing the Ti-6Al-4V alloy powder.

Parameter
Value Laser power (W) 140, 150, 160, 170, 180 Scanning speed (mm/s) 700, 800, 900, 1000, 1100, 1200, 1300, 1400, 1500 Hatch space (mm) 0.06, 0.07, 0.08, 0.09, 0.1 Layer thicknesses (mm) 0.03 Annealing heat treatment was applied to the Ti-6Al-4V samples printed using optimized process parameters. It was reported that the annealing heat treatment for the SLMprinted Ti-6Al-4V alloy was mostly conducted at the temperature of 600-750 °C for 2 h, and the favorable temperature range was 800-900 °C for other heat treatments (except annealing) in the same timespan of 2 h [28]. In addition, the Ti-6Al-4V alloy annealed near to the β phase transition temperature (995 °C) could obtain the most improved mechanical properties [29]. Therefore, the annealing temperatures were set at 750 °C, 850 °C, 950 °C, and 1050 °C respectively. The samples were then held in the furnace for 2 h and then cooled. The procedure of the annealing process is shown in Figure 2.

SLM Process and Heat Treatment
A Dimetal-100 SLM equipment (Laseradd Technology Co., Ltd., Guangzhou, Guangdong, China) was utilized to print the Ti-6Al-4V alloy powder. The process optimization was conducted, and the parameter variables are shown in Table 2. The scanning strategy was bidirectional orthogonal scanning with a scanning starting angle of 145 • . The cubic samples with dimensions of 8 mm × 8 mm × 8 mm and tensile samples were vertically printed, respectively, using various laser power, scanning speed, and hatch space values. All the samples were printed at the same position, i.e., central area of the substrate.
Annealing heat treatment was applied to the Ti-6Al-4V samples printed using optimized process parameters. It was reported that the annealing heat treatment for the SLM-printed Ti-6Al-4V alloy was mostly conducted at the temperature of 600-750 • C for 2 h, and the favorable temperature range was 800-900 • C for other heat treatments (except annealing) in the same timespan of 2 h [28]. In addition, the Ti-6Al-4V alloy annealed near to the β phase transition temperature (995 • C) could obtain the most improved mechanical properties [29]. Therefore, the annealing temperatures were set at 750 • C, 850 • C, 950 • C, and 1050 • C respectively. The samples were then held in the furnace for 2 h and then cooled. The procedure of the annealing process is shown in Figure 2.

Characterizations
After printing, the relative density of the samples was measured to determine the optimal process parameters. The relative density of the printed Ti-6Al-4V samples was

Characterizations
After printing, the relative density of the samples was measured to determine the optimal process parameters. The relative density of the printed Ti-6Al-4V samples was measured by an OHAUS PX124ZH electronic analytical balance (OHAUS Corporation, Parsippany, NJ, USA) according to the Archimedes drainage method. The samples were ground, polished, and etched with Kroll's agent (H 2 O/HF/HNO 3 = 1:3:50 mL) for 20 s for microstructure characterization. The microstructure of the samples was observed with a Leica inverted optical microscope (Leica Microsystems GmbH, Wetzlar, Germany). The phase identification of the printed samples before and after annealing was conducted by a D8 ADVANCE X-ray diffractometer with copper target X-ray (Bruker AXS GmbH, Karlsruhe, Germany), with a scanning speed of 4 • /min, scan angle of 20-80 • , and voltage value of 40 kV. The fracture morphology of the samples was observed by a Quanta 250 scanning electron microscope (FEI Company, Hillsboro, Oregon, USA).

Mechanical Testing
Tensile samples were printed according to ASTM-E8 (65 mm height, 25 mm gauge length, 2 mm thickness, and 5 mm width). Three samples were tested for each scanning speed and annealing temperature. The tensile tests at room temperature were carried out on a CMT5504 electronic universal testing machine (Zhuhai SUST Electrical Equipment Co., Ltd., Zhuhai, Guangdong, China) with an NCS electronic extensometer (NCS Testing Technology Co., Ltd., Beijing, China) and a speed of 0.5 mm/min. Figure 3 shows the relative density of the SLM-printed samples with various scanning speed and laser power values when the hatch space was 0.07 mm, and the layer thickness was 0.03 mm. Under a low laser power (140-150 W), the relative density increased from 98.3% to 99.04% with the scanning speed increasing from 700 mm/s to 900 mm/s, and then decreased to 97.4% when the scanning speed was larger than 900 mm/s (Figure 3a). A low laser power is prone to producing lack of fusion pores [30]. At low scanning speeds, the laser beam can continuously heat the melt pools, resulting in greater laser energy to the Ti-6Al-4V powder particles and more unstable molten pool flow [31]. It is easy to trap gas into the melt pools to form micropores during their solidification [32]. Meanwhile, serious sputtering occurs, and splashed metal particles fall back to the surface of powder bed to form metal spheres. The reason for the decrease in the relative density is that the accumulation of the spheres leads to the generation of inclusions and pores [33].

Relative Density
Comparatively, under a relatively high laser power (160-180 W) (Figure 3b), the relative density could reach the highest values at a scanning speed of 1300 mm/s. A high laser power can result in a large depth of the powder layer penetrated by the laser beam, which improves the fluidity of the melt pool [34]. At high scanning speeds, a large solidification shrinkage of the melt pools tends to occur, resulting in a poor multi-track overlap and large gap between the tracks. The increase in the gap leads to the increase in layer thickness in the track gap after powder spreading. Therefore, the effective energy density is reduced, which promotes the formation of pores and reduces the relative density of the printed samples. The highest relative density of 99.51% could be achieved with a laser power of 170 W and a scanning speed of 1300 mm/s for the SLM-printed Ti-6Al-4V alloy. Meanwhile, serious sputtering occurs, and splashed metal particles fall back to the surface of powder bed to form metal spheres. The reason for the decrease in the relative density is that the accumulation of the spheres leads to the generation of inclusions and pores [33]. Comparatively, under a relatively high laser power (160-180 W) (Figure 3b), the relative density could reach the highest values at a scanning speed of 1300 mm/s. A high laser power can result in a large depth of the powder layer penetrated by the laser beam, which improves the fluidity of the melt pool [34]. At high scanning speeds, a large solidification shrinkage of the melt pools tends to occur, resulting in a poor multi-track overlap and large gap between the tracks. The increase in the gap leads to the increase in layer thickness in the track gap after powder spreading. Therefore, the effective energy density is reduced, which promotes the formation of pores and reduces the relative density of the printed samples. The highest relative density of 99.51% could be achieved with a laser power of 170 W and a scanning speed of 1300 mm/s for the SLM-printed Ti-6Al-4V alloy. Figure 4 shows the variation of the relative density of the SLM-printed samples with hatch space when the laser power was 170 W, the scanning speed was 1300 mm/s, and the layer thickness was 0.03 mm. When the hatch space increased from 0.06 mm to 0.07 mm and from 0.07 mm to 0.1 mm, the relative density of the sample increased from 99.21% to 99.5% and then gradually decreased to 98.5%, respectively. When the laser power and scanning speed were kept constant, the laser input energy was constant, and the melt pool width remained stable. A large hatch space resulted in a quite small overlap rate between adjacent melt pools, which is conducive to the formation of porosity. Comparatively, the decrease in the hatch space increased the overlap rate and reduced the heating time interval between the adjacent melt pools, resulting in a sufficient metal flow within the melt pools and resultant high relative density [35].  Figure 4 shows the variation of the relative density of the SLM-printed samples with hatch space when the laser power was 170 W, the scanning speed was 1300 mm/s, and the layer thickness was 0.03 mm. When the hatch space increased from 0.06 mm to 0.07 mm and from 0.07 mm to 0.1 mm, the relative density of the sample increased from 99.21% to 99.5% and then gradually decreased to 98.5%, respectively. When the laser power and scanning speed were kept constant, the laser input energy was constant, and the melt pool width remained stable. A large hatch space resulted in a quite small overlap rate between adjacent melt pools, which is conducive to the formation of porosity. Comparatively, the decrease in the hatch space increased the overlap rate and reduced the heating time interval between the adjacent melt pools, resulting in a sufficient metal flow within the melt pools and resultant high relative density [35].     Figure 5 exhibits the microstructure of the SLM-printed Ti-6Al-4V samples manufactured at different scanning speeds under a laser power of 170 W, hatch space of 0.07 mm, and layer thickness of 0.03 mm. The microstructure of the samples with different scanning speeds was mainly composed of coarse epitaxial columnar grains that grew along the building direction. In the SLM process, the melt pool temperature is generally higher than that of the β phase generation. The ultrahigh cooling rates (up to 10 6 K/s) suppresses the transformation from the β phase into the α phase, and martensitic transformation occurs to form fine acicular α grains [36]. The primary β grains were filled with fine acicular α martensite that grew toward 45 • upward with the building direction. When the scanning speed ranged from 900 mm/s to 1100 mm/s, the average width of the primary β grain was about 200 µm. However, the further increase in the scanning speed from 1200 mm/s to 1300 mm/s resulted in a decrease in the average width of the primary β grain to 150 µm. In addition, the scanning speed had a great impact on the porosity of the samples, which is consistent with the results shown in Figure 3.  Figure 6 shows the X-ray Diffraction (XRD) pattern of the SLM-printed Ti-6Al-4V samples with different scanning speeds. Both α and α' phases possessed hexagonal closepacked (hcp) structures. The diffraction angle of the strongest peak of the samples was shifted to a large value, as compared to the standard diffraction angle of 40.251°. This peak shift confirmed the formation of the martensite α' phase. The increase in scanning speed led to the larger shift of the α' peak, due to the increase in the cooling rate of the melt pool [37]. The cooling rate increase reduced the β precipitated phase but increased the content of V and Al in the acicular α′ phase, which decreased the lattice size of the α′ phase and the augmentation of its diffraction angle [14]. In the spectrum, the diffraction peak of the β phase was not obvious because of its low content caused by its transformation into the α' phase during the cooling process.  Figure 6 shows the X-ray Diffraction (XRD) pattern of the SLM-printed Ti-6Al-4V samples with different scanning speeds. Both α and α phases possessed hexagonal closepacked (hcp) structures. The diffraction angle of the strongest peak of the samples was shifted to a large value, as compared to the standard diffraction angle of 40.251 • . This peak shift confirmed the formation of the martensite α phase. The increase in scanning speed led to the larger shift of the α peak, due to the increase in the cooling rate of the melt pool [37]. The cooling rate increase reduced the β precipitated phase but increased the content of V and Al in the acicular α phase, which decreased the lattice size of the α phase and the augmentation of its diffraction angle [14]. In the spectrum, the diffraction peak of the β phase was not obvious because of its low content caused by its transformation into the α phase during the cooling process.

Microstructure
shift confirmed the formation of the martensite α' phase. The increase in scanning speed led to the larger shift of the α' peak, due to the increase in the cooling rate of the melt pool [37]. The cooling rate increase reduced the β precipitated phase but increased the content of V and Al in the acicular α′ phase, which decreased the lattice size of the α′ phase and the augmentation of its diffraction angle [14]. In the spectrum, the diffraction peak of the β phase was not obvious because of its low content caused by its transformation into the α' phase during the cooling process.   Figure 7 shows the mechanical properties of the SLM-printed Ti-6Al-4V samples with different scanning speeds. The scanning speed had a significant effect on the elongation but not on the tensile strength. The tensile strength of the samples was in the range of 1200-1265 MPa, and the maximum value could be obtained when the laser power was 170 W and the scanning speed was 900 mm/s. However, the elongation firstly increased from 5.5% to 7.8% and then decreased with an increase in the scanning speed. The maximum elongation could be obtained at a scanning speed of 1300 mm/s. Since the as-printed Ti-6Al-4V samples contained complete martensite structures and fine grains, their tensile strength was much larger than the minimum strength requirements for forged Ti-6Al-4V specified in the standard ASTM F1472-14, but the elongation was lower due to the brittleness of the martensite and large residual stress.

Mechanical Properties
Micromachines 2022, 13, 331 8 of 16 Figure 7 shows the mechanical properties of the SLM-printed Ti-6Al-4V samples with different scanning speeds. The scanning speed had a significant effect on the elongation but not on the tensile strength. The tensile strength of the samples was in the range of 1200-1265 MPa, and the maximum value could be obtained when the laser power was 170 W and the scanning speed was 900 mm/s. However, the elongation firstly increased from 5.5% to 7.8% and then decreased with an increase in the scanning speed. The maximum elongation could be obtained at a scanning speed of 1300 mm/s. Since the as-printed Ti-6Al-4V samples contained complete martensite structures and fine grains, their tensile strength was much larger than the minimum strength requirements for forged Ti-6Al-4V specified in the standard ASTM F1472-14, but the elongation was lower due to the brittleness of the martensite and large residual stress. The scanning speed could influence the amount of the precipitated β phase and the size of the α' martensite in the SLM-printed samples. The β phase distributed at the boundary of the acicular α' martensite phase had higher strength, and the decrease in elongation was mainly due to the dislocation locking by β during tensile stress [38], which hindered the movement of dislocations between the α phases. Therefore, with the increase The scanning speed could influence the amount of the precipitated β phase and the size of the α martensite in the SLM-printed samples. The β phase distributed at the boundary of the acicular α martensite phase had higher strength, and the decrease in elongation was mainly due to the dislocation locking by β during tensile stress [38], which hindered the movement of dislocations between the α phases. Therefore, with the increase in the scanning speed, the β phase decreased and the elongation of the samples increased. On the other hand, the reduction in the slip length of the α phase may have resulted in the increase in elongation [39]. The slip length of the α phase could be approximately equal to the width of the acicular α martensite. Upon increasing the scanning speed, the α martensite was refined, thereby reducing the slip length of the α phase and increasing the elongation. Figure 8 presents the representative tensile fracture morphology of the SLM-printed samples under a scanning speed of 900 mm/s and 1300 mm/s. When the scanning speed was 900 mm/s, the fracture surface was mainly composed of flat cleavage steps, and also contained shallow dimples, indicating brittle fracture. The edge of the sample was exposed with more pore defects and nonmolten powder particles. These defects came from the insufficient laser energy input during the SLM process and were strongly related to the improper scanning speed. During the tensile testing, the pores led to the initiation of cracks, causing premature failure of the sample. Under a scanning speed of 1300 mm/s, the sample showed few internal defects and shearing surfaces. The higher scanning speed reduced the size of the crystal grains, which exhibited higher ductility under tension [12]. the sample showed few internal defects and shearing surfaces. The higher scanning speed reduced the size of the crystal grains, which exhibited higher ductility under tension [12].

Effect of Annealing Temperature on Microstructure
The Ti-6Al-4V alloy samples before annealing were printed with a laser power of 170 W, a scanning speed of 1300 mm/s, and a hatch space of 0.07 mm. Figure 9 presents the XRD pattern of SLM-printed Ti-6Al-4V alloy samples after annealing at different temperatures. The results showed similar α and β diffraction peaks of the samples at different annealing temperatures. However, the diffraction peak of the β phase was weak, indicating the low volume fraction of the β phase. The full width at half maximum (FWHM) of the α/α′ peak (2θ = 40.2°) is listed in Table 3. Compared to the SLM-printed sample, the FWHM of the heat-treated samples significantly decreased, elucidating that the residual stress within the printed samples was significantly eliminated [40].

Effect of Annealing Temperature on Microstructure
The Ti-6Al-4V alloy samples before annealing were printed with a laser power of 170 W, a scanning speed of 1300 mm/s, and a hatch space of 0.07 mm. Figure 9 presents the XRD pattern of SLM-printed Ti-6Al-4V alloy samples after annealing at different temperatures. The results showed similar α and β diffraction peaks of the samples at different annealing temperatures. However, the diffraction peak of the β phase was weak, indicating the low volume fraction of the β phase. The full width at half maximum (FWHM) of the α/α peak (2θ = 40.2 • ) is listed in Table 3. Compared to the SLM-printed sample, the FWHM of the heat-treated samples significantly decreased, elucidating that the residual stress within the printed samples was significantly eliminated [40].  The microstructure of the SLM-printed samples after annealing at different temperatures is shown in Figure 10. When the annealing temperature of 750 °C was applied, the microstructure of the sample changed significantly as compared with that of the asprinted sample (Figure 10a,b). The primary β columnar grains in the annealed sample still existed, while the acicular α' martensite in the columnar grains transformed into a mixed α + β phase. The annealing temperature of 750 °C could only drive the partial decomposition of the α' martensite; thus, the α phase in the structure maintained the acicular shape [41]. At the annealing temperature of 850 °C, there was still epitaxial growth of the β columnar grains, but the β boundary became blurred and disappeared (Figure 10c,d). Compared with the annealed microstructure at 750 °C, the metastable acicular α' martensite phase in the β columnar crystals almost decomposed, and the lath-shaped α phase increased and became coarse. After annealing at 850 °C, the microstructure consisted of the α and β phase, showing a network structure. Studies have shown that the α' phase of Ti-6Al-4V alloy can completely decompose at temperatures above 800 °C [42].
After annealing at 950 °C, β columnar grains disappeared and the lamellar α phase could be observed (Figure 10e,f). The grains were further coarsened, and the distribution of α + β dual phase was more uniform. When the annealing temperature increased to 1050 °C, exceeding the β phase transition temperature, the microstructure was completely β phase in the heat preservation state, and the columnar structure in the as-printed sample could be completely eliminated (Figure 10g,h). Due to the slow cooling rate, the α phase in the β grain gathered to form a lath-shaped structure with the same orientation, and coarse Widmanstätten structures could be obtained [43].  The microstructure of the SLM-printed samples after annealing at different temperatures is shown in Figure 10. When the annealing temperature of 750 • C was applied, the microstructure of the sample changed significantly as compared with that of the as-printed sample (Figure 10a,b). The primary β columnar grains in the annealed sample still existed, while the acicular α martensite in the columnar grains transformed into a mixed α + β phase. The annealing temperature of 750 • C could only drive the partial decomposition of the α martensite; thus, the α phase in the structure maintained the acicular shape [41]. At the annealing temperature of 850 • C, there was still epitaxial growth of the β columnar grains, but the β boundary became blurred and disappeared (Figure 10c,d). Compared with the annealed microstructure at 750 • C, the metastable acicular α martensite phase in the β columnar crystals almost decomposed, and the lath-shaped α phase increased and became coarse. After annealing at 850 • C, the microstructure consisted of the α and β phase, showing a network structure. Studies have shown that the α phase of Ti-6Al-4V alloy can completely decompose at temperatures above 800 • C [42].
After annealing at 950 • C, β columnar grains disappeared and the lamellar α phase could be observed (Figure 10e,f). The grains were further coarsened, and the distribution of α + β dual phase was more uniform. When the annealing temperature increased to 1050 • C, exceeding the β phase transition temperature, the microstructure was completely β phase in the heat preservation state, and the columnar structure in the as-printed sample could be completely eliminated (Figure 10g,h). Due to the slow cooling rate, the α phase in the β grain gathered to form a lath-shaped structure with the same orientation, and coarse Widmanstätten structures could be obtained [43].   Table 4 and Figure 11 show the tensile properties of the SLM-printed samples at different annealing temperatures. It can be seen that the tensile strength decreased gradually with the increase in the annealing temperature. When the annealing temperature of 750 • C was applied, the tensile strength of the sample decreased to 1094 MPa, which is 9% lower than that of the printed sample. The elongation of 7% was similar to that of the printed sample. The maximum elongation of 14% could be obtained at 950 • C, which is 79% higher than that of the as-printed sample.  Table 4 and Figure 11 show the tensile properties of the SLM-printed samples at different annealing temperatures. It can be seen that the tensile strength decreased gradually with the increase in the annealing temperature. When the annealing temperature of 750 °C was applied, the tensile strength of the sample decreased to 1094 MPa, which is 9% lower than that of the printed sample. The elongation of 7% was similar to that of the printed sample. The maximum elongation of 14% could be obtained at 950 °C, which is 79% higher than that of the as-printed sample.  After annealing at 750 °C, some brittle and hard α' martensites decomposed into the α + β phase with relatively high ductility. However, the partial decomposition suppressed the change in the elongation of the sample but reduced its tensile strength. When the annealing temperature exceeded 800 °C, the acicular α' martensite completely decomposed into the α and β phases, which decreased the tensile strength and gradually increased the elongation. When the annealing temperature exceeded the β phase transition temperature, the formed coarse grains and the lath-shaped α phases inside them could hinder the slip of dislocations, causing stress concentration at the interface of the α and β phases and eventually reducing the ductility. Figure 12 shows the fracture morphology of the samples at different annealing temperatures. It can be observed that the fracture morphology of the samples annealed at 750 °C and 850 °C was similar to that of the as-printed samples, including cleavage facets and dimples. However, when the annealing temperature increased to 950 °C, the cleavage fac- After annealing at 750 • C, some brittle and hard α martensites decomposed into the α + β phase with relatively high ductility. However, the partial decomposition suppressed the change in the elongation of the sample but reduced its tensile strength. When the annealing temperature exceeded 800 • C, the acicular α martensite completely decomposed into the α and β phases, which decreased the tensile strength and gradually increased the elongation. When the annealing temperature exceeded the β phase transition temperature, the formed coarse grains and the lath-shaped α phases inside them could hinder the slip of dislocations, causing stress concentration at the interface of the α and β phases and eventually reducing the ductility. Figure 12 shows the fracture morphology of the samples at different annealing temperatures. It can be observed that the fracture morphology of the samples annealed at 750 • C and 850 • C was similar to that of the as-printed samples, including cleavage facets and dimples. However, when the annealing temperature increased to 950 • C, the cleavage facets disappeared and dense dimples with large sizes were formed, indicating a ductile fracture. When the local stress at the phase interface exceeded the interfacial bonding force in the tensile process, micropores occurred and consumed a large amount of strain energy. Micropores lead to dimples in the aggregation process, and denser dimples indicate the better ductility. After annealing at 1050 • C, the dimple size increased, and a small number of tear ridges appeared, which reduced the elongation of the sample.

Effect of Annealing Temperature on Mechanical Properties
Micromachines 2022, 13, 331 13 of 16 ets disappeared and dense dimples with large sizes were formed, indicating a ductile fracture. When the local stress at the phase interface exceeded the interfacial bonding force in the tensile process, micropores occurred and consumed a large amount of strain energy. Micropores lead to dimples in the aggregation process, and denser dimples indicate the better ductility. After annealing at 1050 °C, the dimple size increased, and a small number of tear ridges appeared, which reduced the elongation of the sample.  Table 5 shows the mechanical properties of Ti-6Al-4V samples after different posttreatments as compared to those reported in previous studies. It can be seen that the sample after 950 °C heat treatment exhibited a superior tensile strength and reasonable elongation. Heat treatment temperatures below 950 °C reduced the β phase [23], resulting in a higher tensile strength of the printed alloy. The samples with higher elongations were treated by HIP [26,27], which was more conducive to tailoring the microstructure and reducing the pore defects in the sample.  Table 5 shows the mechanical properties of Ti-6Al-4V samples after different posttreatments as compared to those reported in previous studies. It can be seen that the sample after 950 • C heat treatment exhibited a superior tensile strength and reasonable elongation. Heat treatment temperatures below 950 • C reduced the β phase [23], resulting in a higher tensile strength of the printed alloy. The samples with higher elongations were treated by HIP [26,27], which was more conducive to tailoring the microstructure and reducing the pore defects in the sample.

Conclusions
This work investigated the microstructure and mechanical properties of the SLMprinted Ti-6Al-4V alloy post-treated by annealing. The effects of process parameters on the relative density, microstructure, and mechanical properties of SLM-printed samples were studied. The effects of annealing temperature on microstructure and mechanical properties of the printed samples were further studied. The main findings are presented below.
The relative density of the SLM-printed sample was significantly affected by the scanning speed. In particular, the SLM-printed sample could obtain the highest density of 99.51% with a laser power of 170 W, a scanning speed of 1300 mm/s, a layer thickness of 0.03 mm, and a hatch space of 0.07 mm.
The microstructure of the printed sample was composed of β columnar crystals, which contained a large number of acicular α martensite, resulting in higher strength and lower plasticity of the sample. The width of the β columnar crystals decreased with the increase in the scanning speed, as determined by the decrease in the energy density. The maximum tensile strength of 1265 MPa was achieved at a scanning speed of 900 mm/s, while its elongation could reach the highest value of 7.8% at a scanning speed of 1300 mm/s. The annealing temperature had a significant effect on the microstructure of the sample. After annealing, the acicular α martensite was decomposed into the α + β dual phase. With the increase in the annealing temperature, the tensile strength gradually decreased, while the elongation increased first and then decreased. Annealing at 950 • C could result in the highest elongation of 14%, which is 79% higher than that of the as-printed sample, without a significant reduction in the tensile strength.

Conflicts of Interest:
The authors declare no conflict of interest.